High-strength ultra-thick steel with excellent cryogenic strain aging impact toughness at center zone thereof, and method for manufacturing same

ABSTRACT

An aspect of the present invention is to provide high-strength ultra-thick steel with excellent cryogenic strain aging impact toughness at the center thereof, and a method for manufacturing same. An embodiment of the present invention provides high-strength ultra-thick steel with excellent cryogenic strain aging impact toughness at the center thereof, and a method for manufacturing same, the steel comprising, by wt %, 0.02-0.06% of C, 1.8-2.2% of Mn, 0.7-1.1% of Ni, 0.2-0.5% of Mo, 0.005-0.03% of Nb, 0.005-0.018% of Ti, 80 ppm or less of P, 20 ppm or less of S, and the remainder of Fe and other evitable impurities, wherein the average grain size of grains having a high boundary angle of 15 degrees or greater is 15 μm or less as measured in a range of ⅜t-⅝t in the thickness (t) direction by EBSD.

TECHNICAL FIELD

The present disclosure relates to a high-strength ultra-thick steelmaterial having excellent cryogenic strain aging impact toughness in acenter zone thereof, and a method for manufacturing the same.

BACKGROUND ART

Recently, the development of an ultra-thick high-strength steel materialhas been necessary in the design of structures such as domestic andforeign ships, and when high-strength steel material is used indesigning structures, economic benefits due to reductions of weight ofthe form of the structure may be obtained, and also a thickness of aplate may be reduced, such that ease of processing and welding work maybe secured simultaneously. Also, to improve a transport efficiency ofships, there have been attempts to operate a polar route, and in thiscase, it is expected that demand for a cryogenic toughness guaranteeinghigh-strength and ultra-thick material which may guarantee impacttoughness at −60° C. instead of general steel material guaranteeingimpact toughness at −40° C. could increase.

However, generally, in the case of high-strength steel material, sincesufficient deformation may not occur in an overall structure due to adecrease in the total reduction ratio during the manufacture of aultra-thick material, a structure may become coarse, and particularly,in the case of a center zone, a coarse austenite structure may beformed, such that hardenability may increase and it may be difficult toguarantee impact toughness of the center zone.

Also, when a ship is manufactured, as for a steel material, an originalplate material form may not be used as is and the steel material may beprocessed in the form of a hull through deformation. When suchdeformation is applied to the steel material, impact toughness due tothe deformation may degrade. Also, elements such as carbon and nitrogenmay enter a dislocation created by the transformation over time afterthe transformation, and impact toughness may be further deteriorated dueto the increase in strength. To guarantee this phenomenon, a strainaging impact test to measure impact toughness after heat treatment at250° C. for 1 hour after strain of 5% may be included in test items fora base material when after a steel material is developed and certifiedby each classification society. Therefore, in the case of ultra-thickand high-strength steel material for ships which may guarantee cryogenictoughness, basic impact toughness and also deformation aging impactproperties may need to be guaranteed, but to guarantee deformation agingimpact for even a center zone of an ultra-thick material, it may benecessary to remarkably improve a microstructure of the center zone,which may be problematic.

Accordingly, in a high-strength steel material of 500 MPa or more, itmay be necessary to improve deformation aging impact toughness of acenter zone by controlling impact toughness of ¼t and ½t zone basematerial and also a microstructure of the center zone.

SUMMARY OF INVENTION Technical Problem

An aspect of the present disclosure is to provide a high-strengthultra-thick steel material with excellent cryogenic strain aging impacttoughness in a center zone thereof, and a method for manufacturing thesame.

Solution to Problem

An embodiment of the present disclosure provides a high-strengthultra-thick steel material having excellent cryogenic strain agingimpact toughness in a center zone thereof including, by wt %, 0.02-0.06%of C, 1.8-2.2% of Mn, 0.7-1.1% of Ni, 0.2-0.5% of Mo, 0.005-0.03% of Nb,0.005-0.018% of Ti, 80 ppm or less of P, 20 ppm or less of S, and abalance of Fe and inevitable impurities, wherein an average grain sizeof grains having a high boundary angle of 15 degrees or greater,measured by EBSD, is 15 μm or less in a ⅜t-⅝t zone in a thickness (t)direction.

Another embodiment of the present disclosure provides a method formanufacturing a high-strength ultra-thick steel material havingexcellent cryogenic strain aging impact toughness in a center zonethereof including reheating a steel slab including, by wt %, 0.02-0.06%of C, 1.8-2.2% of Mn, 0.7-1.1% of Ni, 0.2-0.5% of Mo, 0.005-0.03% of Nb,0.005-0.018% of Ti, 80 ppm or less of P, 20 ppm or less of S, and abalance of Fe and inevitable impurities to a temperature of 1000-1080°C.; obtaining a bar by rough-rolling the reheated steel slab at atemperature of 850-1050° C.; obtaining a hot-rolled steel material byfinish-rolling the bar at a temperature of 700-800° C. at a totalreduction ratio of more than 60%; and cooling the hot-rolled steelmaterial to a temperature of 500° C. or less at a cooling rate of 3°C./s or more.

Advantageous Effects of Invention

According to an aspect of the present disclosure, high-strengthultra-thick steel material with excellent cryogenic strain aging impacttoughness in a center zone thereof which may have yield strength of 500MPa or more and a transition temperature of −60° C. or less during astrain aging impact test for a center zone of a thickness, and a methodfor manufacturing the same.

BEST MODE FOR INVENTION

Hereinafter, an embodiment of steel material of the present disclosurewill be described. First, an alloy composition of the present disclosurewill be described. The unit of alloy composition described below may beweight % unless otherwise indicated.

C: 0.02-0.06%

C may be the most important element for securing basic strength in thepresent disclosure, and accordingly, C may need to be included in steelwithin an appropriate range. However, when the content of C exceeds0.06%, a large amount of C may be fixed to dislocation during a strainaging impact test and strength may increase, such that strain agingimpact toughness may decrease, and when the content is less than 0.02%,strength may decrease. Thus, the content of C may be preferably in therange of 0.02-0.06%. A lower limit of C may be more preferably 0.024%,even more preferably 0.028%, and most preferably 0.3%. An upper limit ofC may be more preferably 0.058%, even more preferably 0.054%, and mostpreferably 0.05%.

Mn: 1.8-2.2%

Mn may be a useful element for improving strength through solid solutionstrengthening and hardenability improvement, and accordingly, 1.8% ormore of Mn may need to be added to satisfy yield strength of 500 MPa ormore to be obtained in the present disclosure. However, when the contentexceeds 2.2%, hardenability may excessively increase such that theformation of coarse upper bainite and martensite may be facilitated suchthat strain aging impact toughness of a center zone may greatly degrade.Thus, the Mn content may be in the range of 1.8-2.2% preferably. A lowerlimit of Mn may be more preferably 1.83%, even more preferably 1.86%,and most preferably 1.9%. An upper limit of Mn may be more preferably2.17%, even more preferably 2.14%, and most preferably 2.1%

Ni: 0.7-1.1%

Ni may facilitate cross slip of dislocation and may improve impacttoughness and hardenability, and accordingly, Ni may be an importantelement to improve strength. To improve strain aging impact toughness ofthe center zone in high-strength steel having yield strength of 500 MPaor more, Ni may be added by 0.7% or more. However, when the contentexceeds 1.1%, hardenability may excessively increase, and a large amountof low-temperature transformation phase may be formed, such thattoughness may decrease, and manufacturing costs may increase, which maybe problematic. Accordingly, the Ni content may be preferably in therange of 0.7-1.1%. The Mn content may be preferably in the range of1.8-2.2%. A lower limit of Ni may be more preferably 0.73%, even morepreferably 0.76%, and most preferably 0.8%. An upper limit of Ni may bemore preferably 1.07%, even more preferably 1.03%, and most preferably1%.

Mo: 0.2-0.5%

Mo may be an important element for improving strength by improvinghardenability, and may be an alloying element having less reduction intoughness as compared to strength improvement, preferably, 0.2% or moreof Mo may be added to secure high-strength steel having yield strengthof 500 MPa or more. However, when the content exceeds 0.5%,hardenability may excessively increase, and a large amount oflow-temperature transformation phase may be formed, such that toughnessmay decrease. Therefore, the Mo content may be preferably in the rangeof 0.2-0.5%. A lower limit of Mo may be more preferably 0.23%, even morepreferably 0.26%, and most preferably 0.3%. An upper limit of Mo may bemore preferably 0.48%, even more preferably 0.44%, and most preferably0.4%.

Nb: 0.005-0.03%

Nb may be precipitated in the form of NbC or NbCN and may improvestrength of a base material. Also, Nb dissolved during reheating to ahigh temperature may be very finely precipitated in the form of NbCduring rolling, may prevent recrystallization of austenite, and mayrefine the structure. To obtain the above effect, Nb may be added 0.005%or more preferably. However, when Nb exceeds 0.03%, brittle cracks maybe created in corners of the steel material, and there may be problemsof deterioration of toughness due to formation of excessive precipitateand formation of a large amount of martensite. Therefore, the Nb contentmay be preferably in the range of 0.005-0.03%. A lower limit of Nb maybe more preferably 0.008%, even more preferably 0.011%, and mostpreferably 0.015%. An upper limit of Nb may be more preferably 0.028%,even more preferably 0.026%, and most preferably 0.025%.

Ti: 0.005-0.018%

Ti may be precipitated as TiN during reheating and may prevent growth ofgrains in a base material and a welding heat-affected zone such thatlow-temperature toughness may greatly improve, and Ti may be added by0.005% or more to effectively precipitate TiN. However, when the contentexceeds 0.018%, coarse TiN crystallization may occur such thatlow-temperature toughness may degrade, which may be problematic.Accordingly, the Ti content may be preferably in the range of0.005-0.018%. A lower limit of Ti may be more preferably 0.006%, evenmore preferably 0.008%, and most preferably 0.01%. An upper limit of Timay be more preferably 0.017%, even more preferably 0.016%, and mostpreferably 0.015%.

P: 80 ppm or Less

P may be an element which may cause brittleness at grain boundaries ormay form coarse inclusions, which may lead to brittleness, and toimprove brittle crack propagation resistance, the content thereof may bepreferably limited to 80 ppm or less.

S: 20 ppm or Less

S may be an element which may cause brittleness at grain boundaries ormay form coarse inclusions, which may lead to brittleness. To improvebrittle crack propagation resistance, the content thereof may bepreferably limited to 20 ppm or less.

A remainder of the present disclosure may be iron (Fe). However, in ageneral manufacturing process, inevitable impurities may be inevitablyadded from raw materials or an ambient environment, and thus, impuritiesmay not be excluded. A person skilled in the art of a generalmanufacturing process may be aware of the impurities, and thus, thedescriptions of the impurities may not be provided in the presentdisclosure.

In the steel material of the present disclosure, an average grain sizeof grains having a high boundary angle of 15 degrees or more, measuredby EBSD, in the ⅜t-⅝t zone in a thickness (t) direction may be 15 μm orless, preferably. When the average grain size of grains having a highboundary angle of 15 degrees or more, measured by EBSD, in the ⅜t-⅝tzone in the thickness (t) direction exceeds 15 μm, an effective grainsize due to grain size coarsening may increase, such that an impacttransition temperature may increase, and deformation aging impacttoughness may degrade, which may be problematic.

Meanwhile, a microstructure of the steel material of the presentdisclosure may be a mixed structure including acicular ferrite, granularbainite, upper bainite.

The steel material of the present disclosure may have a thickness of5-90 mm.

The steel material of the present disclosure provided as described abovemay have yield strength of 500 MPa or more. Also, after 5% of strain andperforming heat treatment at 250° C. for 1 hour, a transitiontemperature may be −60° C. or less in the strain aging impact test.

Hereinafter, a method for manufacturing a steel material according to anembodiment of the present disclosure will be described.

First, a steel slab may be reheated to a temperature of 1000-1080° C. Inthe reheating of the steel material of the present disclosure, theheating temperature may be preferably 1000° C. or higher so as to allowcarbonitride of Ti and/or Nb formed during casting to be solid solute.Also, to sufficiently allow carbonitride of Ti and/or Nb to be solidsolute, the heating may be performed to 1030° C. or higher. However,when the reheating is performed to an excessively high temperature,austenite in the center zone may be coarsened, and thus, the reheatingtemperature may be preferably 1080° C. or less, and more preferably1070° C. or less.

The reheated steel slab may be rough-rolled at a temperature of850-1050° C., thereby obtaining a bar. Rough-rolling may be performed tothe reheated slab as above to adjust the shape thereof. Through therough-rolling, destruction of a cast structure such as dendrites formedduring casting and also the effect of reducing the grain size throughthe recrystallization of coarse austenite may be obtained. Meanwhile, torefine the structure by sufficient recrystallization, a total reductionratio during rough-rolling may be 40% or more preferably.

The bar may be finish-rolled at a temperature of 700-800° C. at a totalreduction of more than 60%, thereby obtaining a hot-rolled steelmaterial. In the present disclosure, finish-rolling may be performed topancake an austenite structure of the bar and to obtain dislocation. Thefinish-rolling may be preferably performed at a temperature of 700-800°C. such that the deformation applied to the center zone may bemaintained as much as possible. When the finish-rolling temperature isless than 700° C., ferrite may be precipitated during deformation andboth strength and toughness may be reduced, which may bedisadvantageous. When the temperature exceeds 800° C., the particle sizemay increase, such that impact toughness may deteriorate and sufficientstrength may not be secured, which may be disadvantageous. A lower limitof the finish-rolling temperature may be more preferably 720° C., evenmore preferably 740° C. An upper limit of the finish-rolling temperaturemay be more preferably 780° C., even more preferably 760° C. In thepresent disclosure, to refine the particle size of the center zoneduring the finish-rolling, a total reduction ratio of more than 60% maybe applied preferably. The total reduction ratio during thefinish-rolling may be more preferably 61% or more, and even morepreferably 62%.

The hot-rolled steel material may be cooled to a temperature of 500° C.or less at a cooling rate of 3° C./s or more. When the cooling rate islower than 3° C./s or the cooling stop temperature is more than 500° C.,fine grains may not be properly formed in the present disclosure, suchthat it may be likely that yield strength may be 500 MPa or less.

MODE FOR INVENTION

Hereinafter, the present disclosure will be described in greater detailthrough examples. However, it is necessary to note that the followingexamples are only for describing the present disclosure by examples andnot for limiting the scope of the present disclosure. This is becausethe scope of the present disclosure is determined by the mattersdescribed in the claims and matters reasonably inferred therefrom.

Example

A steel slab having a thickness of 400 mm and an alloy compositionlisted in Table 1 below was prepared, was reheated to a temperature of1040-1070° C., was rough-rolled in a temperature range of 930-1020° C.,thereby obtaining a bar. The bar was finish-rolled under the conditionslisted in Table 2 and a hot-rolled steel material was obtained, and thesteel material was cooled to a temperature of 491-342° C. at a coolingrate of 3.8-5.4° C./sec. A thickness, an average grain size of grainshaving a high boundary angle of 15 degrees or more, measured by EBSD, inthe ⅜t-⅝t zone in a thickness (t) direction, yield strength, and astrain aging impact transition temperature of the center zone (⅜t-⅝t)were measured and listed in Table 3.

In this case, the center zone strain aging impact test was carried bytaking a sample from the center zone of the steel material, performing aheat treatment at 250° C. for 1 hour after 5% of deformation, performingan impact test, and measuring a transition temperature.

TABLE 1 Alloy composition (weight %) Steel type C Mn Ni Mo Nb Ti P (ppm)S (ppm) Inventive steel 1 0.043 1.96 1.05 0.32 0.023 0.017 39 9Inventive steel 2 0.038 2.06 0.87 0.31 0.016 0.009 44 8 Inventive steel3 0.046 1.99 0.79 0.28 0.015 0.013 51 10 Inventive steel 4 0.031 2.131.07 0.43 0.011 0.012 37 7 Inventive steel 5 0.052 1.86 0.94 0.39 0.0210.011 62 13 Comparative steel 1 0.083 2.07 0.86 0.35 0.018 0.013 57 15Comparative steel 2 0.044 2.49 1.06 0.41 0.019 0.011 48 9 Comparativesteel 3 0.016 1.67 0.93 0.39 0.015 0.012 46 13 Comparative steel 4 0.0421.97 0.59 0.36 0.023 0.017 51 11 Comparative steel 5 0.051 2.03 0.940.67 0.019 0.013 38 14 Comparative steel 6 0.039 1.96 0.89 0.33 0.0460.032 38 14

TABLE 2 Finish-Rolling Rough-Rolling Total Cooling Reheating FinishFinish reduction Stop temperature temperature temperature ratio Ratetemperature Classification Steel type (° C.) (° C.) (° C.) (%) (° C./s)(° C.) Inventive example 1 Inventive steel 1 1065 953 735 62 3.7 435Inventive example 2 Inventive steel 2 1072 975 725 61 4.6 488 Inventiveexample 3 Inventive steel 3 1054 892 713 63 5.7 307 Inventive example 4Inventive steel 4 1049 888 749 61 7.9 205 Inventive example 5 Inventivesteel 5 1079 915 755 62 4.4 416 Comparative example 1 Inventive steel 21026 865 769 38 5.1 395 Comparative example 2 Inventive steel 3 1043 903711 49 4.7 407 Comparative example 3 Comparative steel 1 1055 930 736 615.3 453 Comparative example 4 Comparative steel 2 1067 972 744 61 7.1356 Comparative example 5 Comparative steel 3 1037 901 784 63 12.3 415Comparative example 6 Comparative steel 4 1012 859 723 62 3.8 467Comparative example 7 Comparative steel 5 1059 938 733 61 6.5 459Comparative example 8 Comparative steel 6 1038 896 741 62 5.0 437

TABLE 3 Deformation aging impact transition Average grain Yieldtemperature of Thickness size (μm) in strength center zoneClassification (mm) 3/8t-5/8t zone (MPa) (° C.) Inventive 85 13.3 529−71 example 1 Inventive 80 14.3 564 −65 example 2 Inventive 90 12.1 542−7 example 3 Inventive 85 12.8 572 −73 example 4 Inventive 80 14.5 523−64 example 5 Comparative 80 21.2 559 −49 example 1 Comparative 85 18.9556 −51 example 2 Comparative 85 13.9 635 −36 example 3 Comparative 9018.2 693 −31 example 4 Comparative 80 13.5 449 −62 example 5 Comparative80 14.3 508 −44 example 6 Comparative 85 18.7 669 −38 example 7Comparative 80 13.8 609 −37 example 8

In the case of Inventive Examples 1 to 5 satisfying the alloycomposition and manufacturing conditions suggested in the presentdisclosure, the average grain size of grains of the ⅜t-⅝t zone was 15 μmor less, and accordingly, yield strength was 500 MPa or more, and thestrain aging impact transition temperature was −60° C. or less.

In the case of Comparative Examples 1 and 2, the alloy compositionsuggested in the present disclosure was satisfied, but the totalreduction ratio during finish-rolling was low, such that sufficientdeformation was not applied to the center zone, and acicular ferritewhich may greatly affect grain size refinement was not sufficientlyformed, and a large amount of coarse bainite was formed. Accordingly, itis indicated that the grain size as coarsened, the average grain size ofgrains of the ⅜t-⅝t zone exceeded 15 μm, and the strain aging impacttransition temperature of the center zone exceeded −60° C.

In the case of Comparative Example 3, by having a value higher than anupper limit of C suggested in the present disclosure, a large amount ofcoarse bainite phase was formed due to high hardenability, such thatvery high yield strength was exhibited, and although the average grainsize of grains of the ⅜t-⅝t zone was 15 μm or less, a large amount of Cwas fixed to the dislocation during the strain aging impact test, suchthat the strain aging impact transition temperature exceeded −60° C.

In the case of Comparative Example 4, by having a value higher than anupper limit of Mn suggested in the present disclosure, a large amount ofcoarse bainite phase was formed due to high hardenability, such thatvery high yield strength was exhibited, but the average grain size ofgrains of the ⅜t-⅝t zone exceeded 15 μm, and the strain aging impacttransition temperature exceeded −60° C.

In the case of Comparative Example 5, by having a value lower than Alower limit of C and Mn suggested in the present disclosure, a largeamount of soft phase such as polygonal ferrite was formed in the centerzone, and accordingly, yield strength was lower than 500 Mpa.

In the case of Comparative Example 6, by having a value lower than anupper limit of Ni suggested in the present disclosure, although theaverage grain size of grains of the ⅜t-⅝t zone was 15 μm or less, strainaging impact transition temperature exceeded −60° C. due to a decreasein toughness due to the low Ni content.

In the case of Comparative Example 7, by having a higher value than anupper limit of Mo suggested in the present disclosure, a large amount ofcoarse bainite phase was formed due to high hardenability, such thatvery high yield strength was exhibited, but the average grain size ofgrains of the ⅜t-⅝t exceeded 15 μm, and the strain aging impacttransition temperature exceeded −60° C.

In the case of Comparative Example 8, by having a value higher than anupper limit of Ti and Nb suggested in the present disclosure, strengthincreased due to excessive hardenability and the formation ofprecipitate, and the strain aging impact transition temperature exceeded−60° C. due to the decrease in toughness caused by precipitationstrengthening.

1. A high-strength ultra-thick steel material having excellent cryogenicstrain aging impact toughness in a center zone thereof, the steelmaterial comprising: by wt %, 0.02-0.06% of C, 1.8-2.2% of Mn, 0.7-1.1%of Ni, 0.2-0.5% of Mo, 0.005-0.03% of Nb, 0.005-0.018% of Ti, 80 ppm orless of P, 20 ppm or less of S, and a balance of Fe and inevitableimpurities, wherein an average grain size of grains having a highboundary angle of 15 degrees or greater, measured by EBSD, is 15 μm orless in a ⅜t-⅝t zone in a thickness (t) direction.
 2. The steel materialof claim 1, wherein the steel material has a microstructure includingacicular ferrite, granular bainite, and upper bainite.
 3. The steelmaterial of claim 1, wherein the steel material has a thickness of 5-90mm.
 4. The steel material of claim 1, wherein the steel material hasyield strength of 500 MPa or more.
 5. The steel material of claim 1,wherein, after a heat treatment is performed on the steel material at250° C. for 1 hour after deformation of 5%, a transition temperature is−60° C. or less in a strain aging impact test.
 6. A method formanufacturing a high-strength ultra-thick steel material havingexcellent cryogenic strain aging impact toughness in a center zonethereof, the method comprising: reheating a steel slab including, by wt%, 0.02-0.06% of C, 1.8-2.2% of Mn, 0.7-1.1% of Ni, 0.2-0.5% of Mo,0.005-0.03% of Nb, 0.005-0.018% of Ti, 80 ppm or less of P, 20 ppm orless of S, and a balance of Fe and inevitable impurities to atemperature of 1000-1080° C.; obtaining a bar by rough-rolling thereheated steel slab at a temperature of 850-1050° C.; obtaining ahot-rolled steel material by finish-rolling the bar at a temperature of700-800° C. at a total reduction ratio of more than 60%; and cooling thehot-rolled steel material to a temperature of 500° C. or less at acooling rate of 3° C./s or more.
 7. The method of claim 6, wherein atotal reduction ratio during the rough-rolling is 40% or more.